MgZnO based UV detectors

ABSTRACT

Photoconductive devices ( 1,2 ) comprising Mg x Zn 1−x O, that is preferably epitaxially deposited on a substrate ( 21 ), optionally also including a buffer layer ( 22 ), wherein x has a value such that the layer is sensitive to UV light. The a MgZnO device ( 2 ) having predetermined electrical and optical properties and first and second electrodes ( 3 ) deposited on a surface of the device, the second electrode being spaced from the first electrode. A voltage source ( 4 ) is connected across the first and second electrodes to create an electric field within the device. In operation, when the surface of the device upon which the electrodes are deposited is subjected to a photon emission, electron-hole pairs are created within the device and flow within the device because of the electric field.

BACKGROUND OF THE INVENTION

1. Field of the Invention

The present invention relates to the optoelectronic applications ofcertain MgZnO materials and methods for their manufacture. Moreparticularly, the invention relates to the use of the MgZnO materials inultraviolet detection devices.

2. Description of the Prior Art

Photodetectors are broadly defined as devices which respond to incidentelectromagnetic radiation by converting the radiation into electricalenergy, thereby enabling measurement of the intensity of the incidentradiation. A photodetector typically includes some sort ofphotoconductive device and external measurement circuitry.Photodetectors have many practical applications. For instance,photodetectors find use in scientific research (such as in scintillationdetectors), in manufacturing (such as in devices to detect and preventspoilage of products by light contamination), and in safety applications(such as in preventing overexposure of workers to certain radiation).

In many applications it is desirable to detect a particular type oflight, i.e., a certain range of wavelengths. In such an application,light having a wavelength falling outside the range of wavelengths,which is desired to be detected constitutes “noise” to thephotodetector. Noise can cause an erroneous response from thephotodetector. Prior art UV photodetectors have the drawback that theytypically respond to visible light.

Al_(x)Ga_(1−x)N is a compound semiconductor that is ideally suited fordevices in the visible and the ultraviolet parts of the spectrum.GaN—AlGaN based solid state ultraviolet (UV) photo detectors, sensitiveto 200 nanometer (nm) to 365 nm UV radiation, have been actively soughtfor applications including solar-blind UV detection and flame sensing.Due to the direct band gap and availability of Al_(x)Ga_(1−x)N in theentire alloy composition range (0<x<1), GaN—AlGaN based UV detectorshave the advantages of high quantum efficiency, tunability of cut-offwavelengths, and the capability of being fabricated as heterostructures.In recent years, GaN—AlGaN photo conductors and photo diodes of bothSchottky and PIN junctions with good performance have been reported.Aluminum gallium nitride has a direct bandgap which is tunable from 3.4electron volts (or 365 nanometers) at x=0 to 6.2 electron volts (or 200nanometers) at x=1. This makes the material ideally suited for intrinsicultraviolet sensors with high responsivities for wavelengths shorterthan 365 nanometers and essentially no photosensitivity for longerwavelengths. Such sensors can there for detect ultraviolet emissionsfrom flames in the presence of hot backgrounds (such as infraredemission from the hot bricks in a furnace).

Gallium nitride is a wide, direct bandgap semiconductor which has abroad range of potential applications for optoelectronic and highpower/temperature electronic devices. A number of devices have beendemonstrated, including high power, short wavelength (blue, violet)light emitting diodes or LED's, ultraviolet photoconductive detectors,ultraviolet Schottky Barrier Detectors, metal-semiconductor field effecttransistors or MESFETS, high electron mobility transistors or HEMTS andheterojunction bipolar transistors or HBTs. In the past, several groupsof investigators have reported on gallium nitride/aluminum galliumnitride based ultraviolet detectors, including photoconductive, SchottkyBarrier, and p-n-junction ultraviolet detectors based on gallium nitridesingle layers or p-n-junctions. These photo-conductor devices were allof lateral geometry and suffer from several problems. For example, forphotoconductors with gains of 1000, the reported bandwidth has only beenaround 1 kHz, which makes them too slow for many applications. Thisresponse speed problem becomes more severe with the addition of aluminumin the active layer.

The lateral Schottky Barrier devices prepared on p-doped gallium nitridewere also slow because of the large series resistance of the p-typelayer resulting from the lower carrier mobility and in concentrationachievable. Further, for the Schottky devices, back illumination throughthe transparent sapphire substrate side was required. This resulted inpoor quantum efficiencies because of light absorption at the galliumnitride-sapphire interface region where a very high dislocation densityexists.

The detection of ultraviolet (UV) light during daylight conditions is animportant problem for both commercial and military applications. It isdifficult to design a very sensitive detector that can be used in broaddaylight to detect very low levels of UV radiation. The spectraldistribution of radiation from the sun is similar to that of a 6,000degree blackbody radiator. The solar spectral distribution drops offvery sharply below 290 nm due to atmospheric absorption by ozone. As aresult, the earth's surface is essentially dark below 290 nm. Asolar-blind detector can be defined as a device or apparatus that onlyresponds to wavelengths below about 285 nm. Applications for solar-blinddetectors include monitoring lightning events during thunderstorms,detecting ultraviolet laser sources such as excimer lasers or frequencyquadrupled Nd:YAG lasers used as LIDAR sources, and ultraviolettelescope detectors for space platforms.

Many prior art approaches have been proposed to achieve solar blinddetector performance. One approach, described in U.S. Pat. No.4,731,881, uses a series of chemical and color glass filters toaccomplish UV transmission below 285 nm and a sharp cut off, blockingwavelengths longer than 285 nm. The chemical filters consist of anexpensive, single crystal nickel sulfate hexahydrate crystal that hasvery poor thermal and moisture stability, and an organic dye, Cation X,contained in a polyvinylalcohol film to provide UV bandpasscharacteristics. This approach uses a relatively expensive UV sensitivephotomultiplier tube for detection.

Another approach (described in U.S. Pat. No. 4,731,881) uses a rubycrystal with interference filters coated on the two faces. The inputface has a bandpass interference filter that transmits a narrow UV bandat approximately 254 nm and rejects all other wavelengths. The outputface of the ruby crystal is coated with an interference filter thattransmits the ruby fluorescence wavelengths and blocks all otherwavelengths. The performance of this device is limited by the band passand broad band blocking capability of interference filters. A dielectriccoating is limited to a rejection of about 10⁵ outside of the bandpassregion. An out-of-pass-band rejection of approximately 10⁸ is necessaryfor true solar blind detection.

Other approaches (described in U.S. Pat. Nos. 4,241,258 and 5,331,168)use UV sensitive phosphor powders as downconverters. Phosphor powdersare highly scattering and can result in reduced light collectionefficiency.

Visible-blind UV detectors also have great potential in applicationssuch as UV radiometry, flame sensing and missile guidance systems.Currently, photocathodes are the only devices capable of addressingthese applications. Unfortunately, these are bulky, difficult tointegrate with control electronics and in general, require highoperating voltages. Typical prior art devices for achievingvisible-blind UV-detection suffer from either excessively lowtransmission in the UV signal wavelength region or inadequate rejectionof visible light.

Other UV detectors are described in U.S. Pat. Nos. 6,104,074; 5,446,286;5,574,286 and 6,137,123.

SUMMARY OF THE INVENTION

It is an object of the present invention to overcome the above describedand other drawbacks of the prior art.

It is one object of the present invention to provide an ultraviolet (UV)light photodetector which is “blind” to visible light.

It is, therefore, another object of the invention to provide an improved“solar blind” radiation detector.

Another object of the invention is to provide a solar-blind radiationdetector apparatus that detects ultraviolet light at wavelengths below290 nm in the presence of solar illumination.

It is another object of the present invention to provide a UVphotodetector fabricated using MgZnO.

According to an embodiment of the invention a UV detector is providedcomprising a thin film of a material comprising Mg_(x)Zn_(1−x)O, whereinx has a value such that the thin film is sensitive to UV light in thewavelength range of from about 150 nm to about 400 nm.

Another embodiment of the present invention comprises a photoconductivedevice comprising MgZnO, that is preferably epitaxially deposited on asubstrate. The deposited MgZnO may include a buffer layer deposited onthe substrate and a MgZnO film deposited on the buffer layer. Theelectrical and optical properties of the device are controlled byvarying parameters of the deposition process.

A photodetector according to one embodiment of the present inventioncomprises a MgZnO device having predetermined electrical and opticalproperties and first and second electrodes deposited on a surface of thedevice, the second electrode being spaced from the first electrode. Avoltage source may optionally be connected across the first and secondelectrodes to create an electric field within the device. In operation,when the surface of the device upon which the electrodes are depositedis subjected to UV radiation, electron-hole pairs are created within thedevice and flow within the device.

In another embodiment, the present invention comprises a method ofmaking a photodetector having predetermined electrical and opticalproperties. The method comprises fabricating a MgZnO device havingpredetermined electrical and optical properties of the device,depositing a first electrode on a surface of the device, and depositinga second electrode on the surface of the device, the second electrodebeing spaced from the first electrode. In a further step a voltagesource is connected across the first and second electrodes.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 depicts a XRD scan of MgZnO film according to the invention grownon an Al₂O₃(0001) substrate.

FIG. 2 depicts a UV-VIS transmission spectrum of MgZnO thin filmsaccording to the invention deposited on double-side polished sapphire.

FIG. 3( a) depicts an optical micrograph of a MgZnO UV detectoraccording to the invention.

FIG. 3( b) depicts linear I-V curves of the detector of FIG. 3( a).

FIG. 4 depicts the spectral response of a MgZnO UV detector of theinvention under front illumination

FIG. 5 depicts the temporal response of a MgZnO UV detector according tothe invention, and

FIG. 6 depicts a photodetector circuit according to one embodiment ofthe invention.

DETAILED DESCRIPTION OF THE INVENTION

ZnO is a wide band gap (E_(g)=3.3 eV) semiconductor that can be used forultraviolet (UV) photon detection. Its high radiation hardness enablesit to be used in harsh environments. [D. C. Look, D. C. Reynolds, J. W.Hemsky, R. L. Jones, and J. R. Sizelove, Appl. Phys. Lett.75, 811(1999)]. The availability of lattice-matched single-crystal ZnOsubstrates and the relatively low deposition temperatures (100–750° C.)for manufacturing devices composed of ZnO ease the device processing.Owing to its large exciton binding energy (60 meV), ZnO has attractedincreasing attention in recent years for potential low-threshold blue/UVlasers that can be integrated with photodetectors. [A Ohtomo, MKawasaki, Y Sakurai, Y Yoshida, H Komuma, P Yu, Z K Tang, G K Wong, andY Segawa, Mat. Sci. Eng. B54, 24 (1998)]. However, the lack of reliablep-type ZnO hinders any p-n junction based optoelectronic devices. Onlymetal-semiconductor-metal (MSM) structured UV detectors with eitherSchottky or ohmic contacts have been reported. [H. Fabricius, T.Skettrup, and P. Bisgaard, Appl. Optics. 25, 2764 (1986); Y. Liu, C. R.Gorla, S. Liang, N. Emanetoglu, Y. Lu, H. Shen, and M. Wraback, J. Elec.Mat. 29, 69 (2000)]. The cutoff wavelength of these ZnO TV detectors isaround 377 nm. Mg_(x)Zn_(1−x)O which is manufactured by alloying MgOwith ZnO, exhibits the same material advantages as pure ZnO. By varyingthe Mg composition, the band gap Of Mg_(x)Zn_(1−x)O can be tuned from3.3 eV to 7.8 eV for wurtzite and cubic-structured Mg_(x)Zn_(1−x)O,extending the cutoff wavelength from UV-A (320–400 nm) to UV-B (280–320nm) and UV-C (200–280 nm) regions. [A. Ohtomo, M. Kawasaki, 1. Koida, K.Masubuchi, H. Koinuma, Y. Sakurai, Y. Yoshida, 1. Yasuda, and Y. Segawa,Appl. Phys. Left. 72, 2466 (1998); C. W. Teng, J. F. Muth, U. Ozgtir, M.J. Bergmann, H. O. Everitt, A. K. Sharmar, C. Jin, and J. Narayan Appl.Phys. Lett. 76, 979 (2000); A. K. Sharma, C. Jin, A. Kvit, J. Narayan,J. F. Muth, and O. W. Holland, Abstracts of MRS Spring 2000 Meeting,J3.19, 179 (2000)]. Such wide ranges of sensing spectra are expected toenable Mg_(x)Zn_(1−x)O UV detectors to be used in many applications suchas solar UV radiation monitoring, ultra-high temperature flame detectionand airborne missile warning systems. [P. Schreiber, T. Dang, G. Smith,T. Pickenpaugh, P. Gehred, and C. Litton, Proc. SPIE 3629, 230 (1999)].Nevertheless, UV detectors based on Mg_(x)Zn_(1−x)O have not yet beenreported in the prior art. The present invention enables the epitaxialgrowth of high quality Mg_(x)Zn_(1−x)O films on substrates such asc-plane sapphire by pulsed laser deposition, as well as the fabricationand characterization of photoconductive Mg_(x)Zn_(1−x)O UV detectorswith MSM structure.

The pulsed laser deposition system and Mg_(x)Zn_(1−x)O growth processare described with particularity hereinbelow. [See also R. D. Vispute,V. Talyansky, Z. Trajanovic, S. Choopun, M. Downes, R. P. Sharma, T.Venkatesan, M. C. Woods, R. T. Lareau, K. A. Jones, and A. A. Iliadis,Appl. Phys. Left. 70, 2735 (1997)]. In brief, KrF excimer laser pulses(248 nm, 10 Hz, ˜1.7J/cm²) are focused on a Mg_(0.20)Zn_(0.80)O targetwhich was mounted in a high vacuum chamber with a base pressure of˜1×10⁻⁸ torr. The substrate temperature (750° C.), oxygen pressure(1×10⁻⁴ torr) and growth rate (˜0.3 Å/pulse) were previously optimized.[R. D. Vispute, V. Talyansky, S. Choopun, R. P. Shanna, T. Venkatesan,M. He, X. Tang, J. B. Halpern, M. G. Spencer, Y. X. Li, L. G.Salanmace-Riba, A. A. Iliadis, and K. A. Jones, Appl. Phys. Left. 73,348 (1998)]. The film thickness ranged from about 0.1 to about 1 micron.The Mg composition in the film was measured by energy dispersivespectroscopy, and was found to be Mg_(0.34)Zn_(0.66)O. The differencebetween the Mg percentage in the target and in the film was attributedto the high vapor pressure of Zn and the different sticking coefficientsof Zn and Mg. Without any intentional doping, the film was n-type with aresistivity of ˜5×10⁷ Ω-cm, which were measured by the Seeback effectand the transmission line method, respectively. The crystalline qualityof the films was examined using four-circle x-ray diffraction (XRD) andRutherford backscattering spectrometry (RBS) employing the ionchanneling technique. Optical properties were evaluated usingphotoluminescence and ultraviolet-visible (UV-VIS) transmissionspectroscopy (Shimadzü, UV-2501 PC). To characterize the MgZnO UVdetectors, a monochromator with a 150 W xenon lamp and a 1200 lines/mmgrating was used. The spectral response was measured with a low noisecurrent preamplifier (Stanford Research Systems SR571) and a lock-inamplifier (Stanford Research Systems SR530). A semiconductor parameteranalyzer (Hewlett Packard 4155B) was employed for current-voltage (I-V)characterization. A calibrated power meter (Newport 1815-C, 818-UV) anda diaphragm the size of detector mesa to ensure an accurate powermeasurement, which is essential for obtaining a precise responsivityvalue were utilized. For the temporal response measurement, a nitrogengas laser (λ=337.1 nm, τ≦4 ns) was implemented as the excitation source.The power dependence of photoresponse was measured with a continuouswave He—Cd laser (325 nm, 36 mW).

FIG. 1 shows the XRD Θ-2Θ, and Ø scans of MgZnO films grown onAl₂O₃(0001). The appearance of only (0001) peaks in FIG. 1( a) indicatesthat the film is highly c-axis oriented normal to the sapphire (0001)plane. The sharp ω-rocking curve (FWHM=0.14°) for the (0002) peakfurther confirms the good alignment of MgZnO(0001) lattice planes withthe c-plane of sapphire. The in-plane alignment of the film as shown inFIG. 1 c indicates high expitaxial quality. These results confirm thatMg_(0.34)Zn_(0.66)O film is wurtzite-structured and the latticeparameter (c=5.181 Å) is close to that of ZnO (c=5.206 Å). Theappearance of the MgZnO(111) peak at 2Θ=36.95°, though orders ofmagnitude weaker in intensity than that of MgZnO(0002), indicates theonset of phase separation and also the solubility limit of Mg in thewurtzite ZnO lattice. The surface roughness of Mg_(0.34)Zn_(0.66)O filmsis rms˜1.4 nm, as measured by atomic force microscopy. The RBS ionchanneling studies (1.5 MeV He⁺) show a minimum yield (χ_(min)) of ˜5%.The low χ_(min) sharp XRD peaks and smooth morphology indicate goodcrystalline quality of the MgZnO films.

The UV-VIS transmission spectrum of Mg_(0.34)Zn_(0.66)O thin filmsdeposited on double-side polished sapphire is shown in FIG. 2. Thespectra of Mg_(0.18)Zn_(0.82)O and ZnO are also presented for comparisonpurposes. Within the visible region, the average transmittance for allthree films is over 90%. ZnO and Mg_(0.18)Zn_(0.82)O exhibit sharpabsorption edges at 377 nm and 340 nm, respectively, whileMg_(0.34)Zn_(0.66)O at about 308 nm, which is consistent with the peakwavelength of room temperature photoluminescence. From the plot shown inthe inset of FIG. 2, the band gap energy was derived. Clearly, the bandgap increases with the Mg percentage from 3.3 eV for ZnO to 4.05 eV forMg_(0.34)Zn_(0.66)O. The dependence of the Mg_(x)Zn_(1−x)O band gapenergy as a function of Mg content in the films, which was determined byenergy dispersive spectroscopy, is consistent with the results reportedin the prior art.

The MgZnO thin films deposited under optimized condition were utilizedfor UV detector fabrication. Individual detector material was diced andbonded to the TO-9 header for further characterization. FIG. 3 a showsthe optical micrograph of a Mg_(0.34)Zn_(0.66)O UV detector with a sizeof 250×1000 μm². The interdigital metal electrodes, which were definedon ˜1500 Å Cr/Au bilayer by conventional photolithography and ionmilling, are 250 μm long, 5 μm wide and have a pitch of 5 μm. Gold isused due to its excellent conductivity and low tendency to oxidize whendeposited on oxide. However, the as-deposited Au—Mg_(0.34)Zn_(0.66)Ocontacts are Schottky type with a small potential barrier. To ensuregood ohmic contact, a thin layer (˜30 Å) of chromium was used asadhesive followed by a rapid thermal annealing (350° C., 1 min) process.The Cr diffuses into the Mg_(0.34)Zn_(0.66)O and eliminates the contactbarrier. After these treatments, linear I-V curves were obtained asshown in FIG. 3 b. Under 5V bias, the measured average dark current is˜40 nA, which is close to the calculated dark current based on theresistivity of Mg_(0.34)Zn_(0.66)O, indicating that the low dark currentis the direct result of the high resistivity of Mg_(0.34)Zn_(0.66)O. Thelow dark current (I_(dk)) is helpful to enhance the detector's signal tonoise (S/N) ratio since the shot noise, which exceeds the Johnson and1/f noise if the operating frequency is not too low, is proportional toI_(dk):<i _(s) ²>=2q·[1+2(G−1)²](I _(ph) +I _(dk))·B  (1)where q is the electron charge, G is the internal gain, I_(ph) is thephotocurrent, B is the bandwidth.

Upon UV illumination (308 nm 0.1 μW), the photocurrent jumped to 124 μAat 5V bias, indicating a responsivity of ˜1200 A/W. This responsivityvalue is comparable to that of ZnO (400 A/W at 5V bias, 2–16 μminterelectrode spacing) and GaN (2000A/W at 5V bias, 10 μminterelectrode spacing) photoconductive detectors. [See, Fabricus,supra; Ohtomo, supra, and M. AsifKhan, J. N. Kuznia, D. T. Olson, J.M.Van Hove, M. Blasingame, and L. F. Reitz, Appl. Phys. Left. 60, 2917(1992).

The spectral response of a Mg_(0.34)Zn_(0.66)O UV detector under frontillumination is plotted in FIG. 4. The peak response is found at 308 nm,which is in agreement with the absorption edge shown in FIG. 2. The ˜3dB cutoff wavelength is 317 nm, and the visible rejection (R308 nm/R400nm) is more than four orders of magnitude, indicating a high degree ofvisible blindness. For wavelengths beyond the cutoff, the responsivitydrops at a rate of −1.4 dB/nm, followed by a slower decrease, which maybe caused by the Mg_(0.34)Zn_(0.66)O alloy fluctuation. For wavelengthsshorter than 308 nm, the responsivity drops initially followed by anincrease at 210 nm. The drop of the responsivity in this region isattributed to the surface recombination that may be caused by ionmilling. The inset of FIG. 4 shows the responsivity as a function ofbias voltage with 308 nm, 0.1 μW UV illumination. A linear relationshipwas obtained between 0.5V and 5V, indicating no carrier mobilitysaturation or sweep-out effect up to the applied bias.

FIG. 5 shows the temporal response of a Mg_(0.34)Zn_(0.66)O UV detectorwith 3V bias and 50 Ω load. The 10%–90% rise and fall time are 8 ns and˜1.4 μs, respectively. The signal drops to zero at about ˜30 μs asprojected by the dashed line. No persistent photoconductivity wasobserved. The 8 ns rise time is limited by the excitation laser, whichhas a nominal pulse duration FWHM˜4 ns. In order to understand the causeof the much longer fall time, it is necessary to consider variousphysical effects. We fit the decreasing portion of the temporal responsewith an exponential decay curve and found a fast decay component with acharacteristic time τ₁=4.1 ns and a slow decay component with τ₂=3.9 μs.The RC constant may be responsible for the slow decay component.However, the detector capacitance was estimated to be ˜1 pF and wasconfirmed by the actual capacitance-voltage (C-V) measurement. Even withthe inclusion of parasitic capacitance and resistance from the wiringand contacts, the RC product still gives a much shorter time constantthan the observed decay. Thus the RC limit is ruled out. The transittime limit can also be ruled out: At a bias of V_(b)=3V, the transittime T_(tr)=s²/(μ_(n)·V_(b)) is on the order of ˜1 ns, where s is theinterelectrode spacing and μ_(n) is the electron mobility (estimated tobe 50˜100 cm²/Vs)[S. Choopun, R. D. Vispute, W. Noch, A. Balsamo, R. P.Shanna, T. Venkatesan, A. A. Iliadis, and D. C. Look, Appl. Phys. Left.75, 3947 (1999)]. The T_(tr) may contribute to the 8 ns rise time buthas a negligible effect on the fall time. Another factor that affectedthe temporal response is the excess lifetime of trapped carriers,especially the trapped holes in n-type semiconductors. If this is thecase, one can expect a large internal gain and a high currentresponsivity.

The latter can be expressed [M. Razeghi and A. Rogalski, J. Appl. Phys.79, 7433 (1996)] as:R=qλ/hc[η _(in)(1−r)]τ_(p)·(μ_(n)+μ_(p))V _(b) /s ²  (2)where q is the electron charge, λ is the light wavelength, h is thePlank constant, η_(in) is the internal quantum efficiency with a valueclose to unity, μ_(n) and μ_(p), are electron and hole mobility,respectively, and r is the reflectivity which can be derived from therefractive index of Mg_(x)Zn_(1−x)O reported by Teng et al, supra.Assuming the mean lifetime of holes (τ_(p)) is on the same order ofmagnitude as the 10%–90% fall time, the responsivity was estimated to beR˜10³ A/W, which is consistent with the measured responsivity shown inFIG. 3 and FIG. 4. Further evidence comes from the distribution of holelifetime. Assuming the trapped holes have a mean lifetime of τ_(p)˜1.4μs, the plot of hole lifetime distribution P(t)=1/τ_(p) exp(−t/τ_(p))fits well with the temporal response decay curve. It is believed thatthe origin of trap states may relate to the surface damage during thedevice processing, or the interface states associated with thewurtzite-cubic phase separation and alloy fluctuation. It is worthwhileto note that the product of gain and bandwidthG·B=(μ_(n)+μ_(p))·V_(b)/2πs² is a constant on the order of ˜10⁹ Hz.Trading off between the gain and the bandwidth, the Mg_(0.34)Zn_(0.66)OUV detectors are useful for many practical applications that needrelative high gain and moderate bandwidth.

Thus, according to the present invention, there are provided visibleblind UV detectors based on Mg_(x)Zn_(1−x)O thin films having a highresponsivity of 1200 A/W, a fast response of 8 ns rise time, and 1.4 μsfall time. The detector shows peak responsivity at 308 nm and ˜3 dBcutoff at 317 nm. Visible rejection is more than four orders ofmagnitude.

Preferably, Mg_(x)Zn_(1−x)O thin films wherein x has a value betweenabout 0 and about 1 are suitable for the practice of the invention. Mostpreferably the UV thin film detector material comprisesMg_(0.34)Zn_(0.66)O. Thin films of Mg_(x)Zn_(1−x)O are suitablydeposited, preferably epitaxially by the pulsed laser ablationdeposition (PLAD) techniques described hereinbelow, on suitablesubstrates such as Al₂O₃, MgO, ZnO, buffered silicon, SiC, GaN,GaN/Al₂O₃, and the like to a thickness of from about 0.1 to about 1micron.

One embodiment of a photodetector 1, according to the present inventionis depicted in FIG. 6. Photodetector 1 includes a device 2 havinginterdigitated electrodes 3 formed thereon connected to externalcircuitry. The external circuitry comprises a biasing source 4, and ameasurement instrument, such as a current meter 5. Biasing source 4 maybe any suitable voltage source. In a preferred embodiment the source isa battery operating within the range of 5–25 volts. Current meter 5 maybe any suitably sensitive current meter. In one embodiment a KiethleyModel 614 is used.

In operation, photons (in the form of incident light) strike the surfaceof device 2. Photons of suitable wavelength are absorbed and give riseto electron hole pairs (one for every photon absorbed) within device 2.The electrical conductivity of device 2 increases in proportion to thephoton flux (number of photons per second). An external electric field,generated by application of the bias voltage from source 4, causes theelectrons and holes to be transported within the device, thereby givingrise to a current in the external circuitry which is measurable bycurrent meter 5. It will be understood by those skilled in the art,however, that a device embodying the detector of the invention willoperate without an applied voltage inasmuch as an electrical signaloutput is generated by the electron hole pairs created by the absorbedphotons.

Device 2 consists of a thin film of Mg_(x)Zn_(1−x)O formed on asubstrate. Depending on the material used for the substrate, a bufferlayer may be used between the substrate and the thin film. That is,where the lattice mismatch, as defined by the lattice constants and thecrystal structure of the substrate and the single crystal layer islarge, a buffer layer may be used in order to form the single crystallayer on the substrate with reduced defect density. Where the thermalcoefficients of the respective materials are mismatched, a buffermaterial may also be required. Where the lattice constant and thecrystal structure of the substrate and sin crystal layer are about thesame or the thermal coefficients are about the same, a is not necessary.

According to one embodiment, device 2 comprises three layers including abase substrate 21, a buffer layer 22 and a film 23. Base substrate 21is, e.g., silicon. Buffer layer 22 is deposited thereon. Thin film 23 isdeposited on buffer layer 22. Thereafter, the surface of device 2 may becoated with an antireflective coating (not shown).

Following is a description of the epitaxial growth ofcubic-Mg_(x)Zn_(1−x)O thin films on Si(100) using a pulsed laserdeposition technique. To overcome the large lattice and thermal mismatchbetween Mg_(x)Zn_(1−x)O and Si(100), various materials were tested asbuffer layers. The epitaxial growth of cubic-Mg_(x)Zn_(1−x)O on Si(100)was realized with double buffer layers of SrTiO₃/SrO, SrTiO₃/TiN orBiTiO₁₂/Y-stablized-ZrO₂. The epitaxial relationship found to beMg_(x)Zn_(1−x)O(100)// buffer(100)//Si(100), Mg_(x)Zn_(1−x)O[100]//Si[100] for all buffers except for SrTiO₃/SrO, which has a 45°in-plane lattice rotation with respect to Si. The good crystallinequality of the Mg_(x)Zn_(1−x)O epilayers was confirmed by the narrowx-ray diffraction lines, low Rutherford backscattering ion channelingyield and smooth surface morphology observed with atomic forcemicroscope. The Mg and Zn composition in the Mg_(x)Zn_(1−x)O epilayerwas measured using energy dispersive spectroscopy. The band gap energiesof Mg_(x)Zn_(1−x)O films was derived from the ultraviolet-visibletransmission spectra and were found to increase monotacally with thedeposition temperature. The relationship of Mg_(x)Zn_(1−x)O band gapenergy, Zn/Mg composition and the deposition temperature is alsodiscussed herein below.

ZnO and Mg_(x)Zn_(1−x)O have been subjects of intense scientificresearch as wide band gap optoelectronic (OE) materials, due to theever-increasing demands for blue and ultraviolet (UV) photon emittersand detectors in many technical areas. Particularly, the large excitonbinding energy of ZnO (60 meV) enables high exciton population andpotentially low lasing threshold when used in current injection lasers.The availability of lattice matched single-crystal substrates (ZnO, MgOetc.), the tunable band gap of Mg_(x)Zn_(1−x)O (3.3eV to 7.98eV,depending on Mg fraction) and the relatively low thin film growthtemperatures (100–750° C.) ease the device fabrication process. The highchemical inertness hardness enables the ZnO and Mg_(x)Zn_(1−x)O baseddevices to be used in harsh environment. In recent years, roomtemperature stimulated photon emission have been observed by severalinvestigators. Previous studies have shown that ZnO and Mg_(x)Zn_(1−x)Owith low Mg fraction have wurtzite structure, while higher Mgcompositions lead to cubic Mg_(x)Zn_(1−x)O (c-MgZnO) with latticeconstants close to that of MgO. The c-MgZnO has a band gap range thatcovers the solar-blind window (about 240–280 nm), thus rendering itideal for applications such as atmosphere ozone monitoring and missilewarning systems.

To date, most ZnO and Mg_(x)Zn_(1−x)O based OE devices are grown onsapphire and used as discrete units. The integration of ZnO andMg_(x)Zn_(1−x)O-based OE devices with Si-based electronic supportingcircuits on a single chip is of great importance for practicalapplications such as vertical cavity surface emitting laser (VCSEL)array, UV camera and optical computation. Since Si-based VLSI techniquesare well established and their reliability and economy have beenacknowledged for decades, it is technically important and commerciallybeneficial to epitaxially grow ZnO and Mg_(x)Zn_(1−x)O on Si, which isthe bridge toward Si-based optoelectronic integrated circuits (OEIC).However, the epitaxial growth of c-MgZnO on Si(100) is also a challengebecause of the large lattice mismatch (−22.4%) between cMgZnO andSi(100). The thermal mismatch can also not be neglected. Moreover, theeasy oxidation of Si surfaces to yield amorphous SiO₂ layers thatprevent any epitaxial growth, especially since Si can take oxygen fromeither the ambient gas inside the growth chamber or react with thedeposited oxide film at elevated temperature, renders the growth of suchfilms even more problematic. Thus, an appropriate buffer layer must beused and the right growth process must be followed to obtain highcrystalline quality c-MgZnO epilayers.

EXAMPLES

The c-MgZnO films described below as well as the various buffer layerswere grown using PLD technique. In brief, the KrF excimer laser pulses(λ=248 nm, τ=30 ns) were focused on a selected PLD target mounted on amulti-target carousel inside a high vacuum chamber with a base pressureof ˜1×10⁻⁸ torr. The ablated target materials were deposited on heatedsubstrates located ˜7 cm away. The laser energy fluence was set between1.5 J/cm² to 2.5 J/cm², which was previously optimized. Laser pulserepetition rates were varied between 1 to 10 Hz. The substratetemperatures and the oxygen pressures inside the chamber were adjustedfrom room temperature to 900° C. and from vacuum to 1×10⁻² torr,respectively.

TABLE 1 Crystal Structures, lattice constants, thermal expansioncoefficients of c- MgZnO, Si and various buffer materials The value ofbuffer-to-Si and buffer-to-c- MgZnO lattice mismatch is calculated ThePLD conditions used in the examples are also given in the table. ThermalLattice Lattice expansion mismatch mismatch to Crystal Latticecoefficient- to Si cMgZnO PLD Materials structure (Å) ×10⁻⁶/° C. (%) (%)conditions c-MgZnO cubic 4.22 ~9 −22.42 — 30–850'° C. 10⁻⁷–10⁻² torr O₂SrTiO_(3*) cubic 3.905 9 +1.6 −7.8 600–800° C. (STO) 10⁻⁴–10⁻² torr O₂SrO cubic — — — — 800° C. vacuum TiN cubic 4.242 9 −21.89 — 800–900° C.vacuum Bi₄Ti₃O₁₂ Ortho- a = 5.41 11 −0.38 — 600–800° C. lci'− (BTO)rhombic b = 5.49 10⁻¹–10⁻⁴ torr O₂ c = 32.8 ZrO₂ cubic 5.139 10 −5.37−18 700–800° C. with 10% Y 10⁻⁵ torr O₂ (YSZ) CeO₂ cubic 5.411 11 −0.36— 700–800° C. 10⁻⁴ torr O₂ Si diamond  a = 5.431 4 — — — *after a 45Uin-plane lattice rotationTable 1 lists the crystal structures, lattice constants and thermalexpansion coefficients of c-MgZnO, Si and various buffer materials usedin the examples. The lattice mismatches of buffer layers to Si andc-MgZnO were calculated. The PLD growth parameters for each material aregiven in the table. All PLD targets arc commercially available exceptthe Mg_(x)Zn_(1−x)O target, which was made by mixing 20% to 50%/mol ofMgO power (4N purity) with ZnO powder (4.5N purity) followed by ahydraulic press (300 Mpa/cm²) and 3 MW microwave sintering (1300° C., 2hr). The target density was over 90% of theoretical.

Except for depositing YSZ, the Si(100) wafer was cleaned by conventionalHF containing solution to remove oxide before loading into the PLDchamber. The H-terminated Si was degassed at 200° C. in vacuum andraised to 850° C. to desorb residual SiO_(x). In the case of depositedYSZ, native silicon oxide was kept on the Si wafer. The thicknesses ofthe c-MgZnO films are 2000–5000 Å, while the buffer layers were keptbelow 500 Å. The crystalline quality of the films was examined usingfour-circle x-ray diffraction. (XRD), and Rutherford backscatteringspectrometry (RBS) with ion channeling technique (1.5 Mev⁴He⁺). Surfacemorphology of the films was observed using atomic force microscope (AFM,Digital Instruments Inc) and scanning electronic microscope (SEM, FEIF1B620). The exact Mg and Zn compositions in the c-MgZnO epilayers weremeasured using energy dispersive x-ray spectrometer (EDS, JEOL 8900).The depth profile of each element in the films was sensed by secondaryion mass spectroscopy (SIMS). Ultraviolet-visible (UV-VIS) transmissionspectroscopy (Shimadzu, UV-2501 PC) was employed to characterize theoptical properties of the c-MgZnO films.

Example 1

FIG. 1 is an XRD Θ-2Θ scan of (a) c-MgZnO/Si(100), (b)c-MgZnO/STO/SrO/Si(100), (c) c-MgZnO/STO/TiN/Si(100) (d) c-MgZnO/BTO/YSZ/Si(100). Thec-MgZnO films were grown at 600° C. with oxygen partial pressure of1×10⁻⁶ torr. The film thickness is 2000–5000 Å. The plot is insemi-logarithm scale. FIG. 1 is the XRD Θ-2Θplot of c-MgZnO films grownon Si(100) with various buffer layers. The result of c-MgZnO films growndirectly on Si(100) without any buffer is also provided as a reference.As can be seen from FIG. 1(a) that without a buffer layer, only a broadMgZnO(111) peak appears at 2Θ=36.8° (FIG. 1-a). The intensity of thispeak is more than four orders of magnitudes lower than that of Si(400),indicating that the film is essentially amorphous. Due to the largelattice mismatch, the epitaxial growth of c-MgZnO on Si(100) can only beobtained when proper buffers are applied. FIGS. 1( b)(c)(d) show thatthe crystalline quality of the c-MgZnO films was improved significantlywhen a buffer was inserted. The high intensity sharp MgZnO (200) peak at2Θ=42.9° indicates that the c-MgZnO is c-axis oriented. The latticeconstant of c-MgZnO is calculated as a=4.22 Å, which is very close tothat of MgO (a=4.22 Å). Note that the buffer layers STO, BTO, TiN andYSZ are also highly c-axis oriented as indicated by the presence only of(001) family peaks, implying that the cubic-on-cubic epitaxial growthhas been realized.

Example 2

FIG. 2 is an XRD Φ scan of c-MgZnO/STO/SrO/Si (b) c-MgZnO/BTO/YSZ/Si.The in-plane alignment of c-MgZnO, buffer layer and Si is indicated bythe position of diffraction peaks. The epitaxial growth of c-MgZnO on Si(100) is further confirmed by the XRD-scan Φ shown in FIG. 2. Forc-MgZnO grown on all three sets of buffer layers, i.e. STO/SrO, STO/TiN,BTO/YSZ, the Φ scans show 4-fold azimuth symmetry, indicating the cubiccrystal structure of c-MgZnO. The average full width at half maximum(FWHM) of c-MgZnO is 0.5° for c-MgZnO with STO/SrO buffer, 0.7° forc-MgZnO with BTO/YSZ buffer, compared with the 0.4° of Si(100). The linewidth of c-MgZnO grown on STO/TiN/Si(100) is slightly larger (˜1°) dueto the poor interface between TiN and Si.

Except for STO/SrO, the maximum Φ-scan diffraction of c-MgZnO is at thesame angle as that of buffer layer and Si substrate, indicating thein-plane alignment is c-MgZnO [100]//buffer [100]//Si[100]. For STO/SrObuffer, however, the four STO(111) peaks are 45° away from that ofSi(111) peaks, implying a 45° in-plane rotation of STO unit cell withrespect to Si. This lattice rotation of STO significantly reduces thelattice mismatch from the −28.2% without rotation to +1.6% with 45°rotation. It is interesting to note that even though the latticemismatch is as large as 22.9%, no lattice rotation was observed betweenTiN and Si. SIMS and TEM results (not shown) indicate that TiN hasdiffused into Si at elevated temperature thus partly releasing thetension. SEM pictures also show the cracking of TiN films along [100]and [100] directions.

Combining the results depicted in FIGS. 1 and 2, the epitaxialrelationship can be described as c-MgZnO(100)//buffer(100)//Si(100) forall three sets of buffer layers, and MgZnO[100]//buffer [100]//Si[100]for STO/TiN and BTO/YSZ, respectively. The STO/SrO buffer rotates 45°from that of the Si lattice with in-plane alignment of c-MgZnO[100]//STO [100]//Si[110].

Example 3

FIG. 3 depicts full width at half maximum (FWHM) of c-MgZnO(200) rockingas curves functions of (a) PLD deposition temperature and (b) oxygenpressure. The c-MgZnO films are grown on STO/SrO/Si(100) (*),BTO/YSZ/Si(100) (**) and STO/TiN/Si(100) ( ). To find the optimal PLDgrowth conditions, the XRD line widths of c-MgZnO grown at differentconditions were measured, which serve as an indicator of c-MgZnOcrystalline quality. Shown in FIG. 3 is the FWHM of co-rocking curves ofc-MgZnO as functions of substrate temperature and oxygen partialpressure. As shown in FIG. 3( a), for substrate temperatures from 30°C., to 600° C., the FWHM of ω-rocking curves decrease sharply from 2.5°to 0.23°, indicating the significant improvement of crystalline quality.Above 600° C., the narrowing of a XRD line width is not significant. At800° C., a broadening of XRD rocking curve to 0.28° was observed, whichmay be caused by the change of chamber base pressure due to thedegassing of the substrate heater. Thus 400° C. to 600° C. is the properwindow of PLD growth.

The c-MgZnO crystalline film can be also improved by lowering the oxygenpartial pressure during the deposition. FIG. 3( b) shows that at fixedsubstrate temperature of 600° C., the rocking curve line width dropsfrom 0.64° at an oxygen pressure of 1×10⁻³ torr to 0.23° if the oxygenpressure drops to 1×10⁻⁶ torr. The results indicate that lower oxygenpartial pressures are beneficial to c-MgZnO growth.

However, as in other PLD processes, the oxygen partial pressure alsoaffects the c-MgZnO film deposition rate. With 10⁻² torr of oxygeninside the chamber, the deposition rate was around 0.5 Å/pulse. Thisnumber drops to 0.3 Å/pulse for oxygen pressure equal to 10⁻⁴ torr, andfurther drops to 0.2 Å/pulse when the oxygen partial pressure is 10⁻⁶torr. Trading off between crystalline quality and deposition rate, theoptimal O₂pressure regime was found to be around 10⁻⁵–10⁻⁶torr.

Other PLD parameters such as laser energy fluence and pulse repetitionrate may also affect c-MgZnO film quality. Within the optimized laserenergy fluence (1.5˜2.5 J/cm²), the variation of XRD line width isnegligible. Beyond 3 J/cm² particles were observed on the c-MgZnO filmsand thus should be avoided. Up to 10 Hz, the effect of laser pulserepetition rate to the c-MgZnO film crystalline quality is not obvious.However, low pulse repetition rate gives low deposition rate and, hence,is not practicable.

Example 4

FIG. 4 depicts RBS random and channeling spectra of c-MgZnO grown onSrTiO₃ buffered Si(100) at temperature (a) 600° C. and (b) 800° C. Thec-MgZnO films were further characterized using RBS with 1.5 MeV He⁺ions.The minimum ion channeling yield χ_(min) which serves as anotherindicator of the film's crystalline quality, was found to also depend onthe growth temperature. Shown in FIG. 4 are the RBS random and alignedbackscattering spectra of c-MgZnO grown at 600° C. and 800° C.,respectively. For the films grown at 600° C., the minimum ion channelingyield is χ_(min)=5% for Zn at the surface. The χ_(min) for Sr atinterface is 12% due to the 8% lattice mismatch between the c-MgZnO toplayer and the STO buffer. This value goes up to 20% if the c-MgZnO filmsare grown at 800° C., which is in accordance with the broadening of theXRD line width shown in FIG. 4. The shape and position of the RBS peakscorresponding to Sr and Zn reveal the reason for the film crystallinedegradation with the increase of temperature. In FIG. 4( a), clear Srand Zn peaks can be resolved, indicating the sharp MgZnO/SrTiO₃interface. In FIG. 4( b), however, no distinct Sr and Zn peaks can befound, even though the Sr, which is part of the STO buffer layer, wasbelow the c-MgZnO top layer. The merging of Sr and Zn spectra indicatethat Sr has diffused into c-MgZnO film at this temperature. Thus roughinterface and higher χ_(min) is expected.

On the low temperature side, though no buffer layer element diffusionwas observed, the χ_(min) also quickly goes up until there is no obviousion channeling at Ts <200° C., indicating the degradation of c-MgZnOcrystalline quality.

Example 5

FIG. 5 sets forth an atomic force microscope picture of c-MgZnO grown onSrTiO₃ buffered Si(100) at temperatures of (a) 600° C. and (b) 800° C.The morphology of c-MgZnO thin films was evaluated using atomic forcemicroscopy. FIGS. 5( a) and (b) are the AFM images of c-MgZnO grown onSrTiO₃ buffered Si at 600° C. and 800° C., respectively. The films grownat 600° C. have a very smooth surface with root mean square (RMS)surface roughness of only 3.57 Å. No obvious grains were observed. Withthe increase of substrate temperature, the films turn rougher, asindicated by the 21.29 Å surface RMS for the film deposited at 800° C. Asimilar trend of surface roughness varying with the depositiontemperature was found for the c-MgZnO films grown on two other sets ofbuffered layers, i.e. BTO/YSZ or STO/TiN. The coincidence of minimum XRDline width shown in FIG. 3, lowest RBS ion channeling in FIG. 4 and thesmooth morphology shown in FIG. 5( a) for the films grown at 600° C.indicate this temperature is within the optimal growth window forc-MgZnO.

Example 6

FIG. 6 shows UV-VIS transmission spectra of c-MgZnO grown MgO(100) at(a) room temperature (b) 200° C. (c) 400° C. (d) 600° C. and (e) 800° C.The results of ZnO (f) and MgO (g) deposited on c-plane sapphiresubstrates were also given for references. The inset is the plot of(αhv)² versus hv-that gives the estimated band gap energy of c-MgZnOdeposited at different temperatures. To investigate the opticalproperties of c-MgZnO films, the c-MgZnO was deposited on MgO(100)substrates. Shown in FIG. 6 are the UV-VIS transmission spectra ofc-MgZnO films deposited on double side polished MgO(100) substrates atdifferent substrate temperatures. The spectra of MgO and ZnO grown onc-plane sapphire substrates are also provided as references. Within thevisible region, the average transmittance of c-MgZnO films is over 90%,indicating that c-MgZnO is highly transparent to visible light.Interference caused oscillation indicates that the films are flat. Fromthe oscillation period, the film thickness was estimated to be 2000–3500Å, which is in agreement with the film thickness measured with RBS andstylus.

It is worthwhile to note that the position of the exciton absorptionedge shifts significantly to the low wavelength side with the increaseof substrate temperature, implying the expansion of c-MgZnO's band gapenergy. This band gap shift is more clear with the plot of (αhv)² versushv shown as the inset of FIG. 6, which gives the band gap energy at theintersection with the horizontal axis. When grown at RT, the c-MgZnOfilm show an absorption edge at 280 nm, corresponding to a band gapenergy of 4.6 eV. The band gap increases nearly linearly with the growthtemperature. At 800° C., the absorption edge shifts to below 200 nm. Theband gap of this film was estimated to be about 6.8 eV.

It is also worthwhile to note that the slope of the absorption edge,which indicates the crystalline quality of the c-MgZnO films, changeswith the deposition temperature. The films deposited at hightemperatures exhibit sharp absorption edges, while film deposited atroom temperature has much flatter absorption edge. The change ofabsorption edge slope with temperature is consistent with the decreaseof XRD line width shown in FIG. 3.

Example 7

FIG. 7 shows the energy dispersive spectroscopy of c-MgZnO filmsdeposited at 600° C. The inset shows the variation of Mg composition inthe films as a function of substrate temperature. The bandgap of c-MgZnOdepends on the Mg and Zn composition in the c-MgZnO films. Hence knowingthe composition of the Mg and Zn in the films is the key tounderstanding the blue shift of the absorption edge shown in FIG. 7.FIG. 7 is the energy dispersive spectroscopy of c-MgZnO films depositedat T_(s)=600° C. on STO/SrO/Si(100). The Zn is marked by the Kα and Lαpeak at 8.638 keV and 1.009 keV, respectively. A small Kα peak at anenergy of about 9 keV is also visible. The Mg is labeled by the Kα peakat 1.254 keV. Within the detection limit of EDS (˜0.5 wt %), nodetectable impurities were found. The Mg and Zn compositions werecalculated and were found to be Mg_(0.78)Zn_(0.21)O for this film. Shownin the inset is the variation of Mg composition in the films as afunction of substrate temperature. The curve is the polynomial fittingof the data point in a least squares sense. As can be seen from the plotthe Mg composition increases almost linearly with temperatures up to600° C. The reason for this increase is discussed herein below.

(a) The selection of buffer materials—To select a proper buffer layer,there are three major concerns; (1) lattice. Thermal match to c-MgZnO ifthe c-MgZnO grows thereon; (2) lattice/thermal match to Si (100) if itgrows thereon, and (3) capability of quenching the formation of SiO_(x)if it grows on Si(100). The first two conditions require the bufferlayer to have a cubic crystal structure and a proper lattice/thermalexpansion constant that can bridge the large lattice mismatch between Siand c-MgZnO. Satisfying the third requirement is essential for theepitaxial growth of any thin films on Si.

SrTiO₃ has a cubic perovskite structure and is well recognized as a highdielectric constant (˜300) material for non-volatile memories and abuffer layer for high-T_(c) superconductors. Its lattice mismatch to Si(100) and c-MgZnO is +1.6% (after a 45° lattice rotation) and −9%respectively. The thermal mismatch is however, tolerable. Testingresults (not shown) of growing c-Mg_(x)Zn_(1−x)O on single crystalSrTiO₃ (100) indicates a high quality epitaxial film, thus renderingSrTiO₃ one of the best choice for buffering c-Mg_(x)Zn_(1−x)O on Si.

Nevertheless, when directly grown on Si(100), SrTiO₃ films usually showa mixture of (100) and (110) planes due to the presence of SiO_(x).Since the oxidation of Si surface is inevitable in the oxide growthprocess, it is necessary to find a second buffer between Si and SrTiO₃that has a larger binding energy with oxygen than the Si—O bond.Accordingly, SrO and TiN can be used to improve the SrTiO₃/Siinterfaces. Formation enthalpy for SrO and Si₃N₄ are −592 kJ/mol and−743.5 kJ/mol, respectively, both being compatible in magnitude withthat of SiO₂'s −910.7 kJ/mol value. Consequently the SrO and TiN quench,to a certain degree, the formation of amorphous SiO_(x). The SrTiO₃films grown on SrO or TiN buffered Si(100) show improved epitaxialquality.

Yttrium stabilized ZrO₂ (YSZ) is another widely used buffer materialthat grows epitaxially on Si(100) even without intentional removal ofnative silicon oxide from the Si substrate. The lattice mismatch betweenYSZ and Si is only −5.37%. The formation enthalpy of ZrO₂ (−100.6kJ/mol) is larger than that of SiO. The combination of these two effectsleads to high quality YSZ epilayers on Si(100) with the epitaxialrelationship of YSZ(100)//Si(100) and YSZ[100]//Si[100].

However, due mainly to the large lattice mismatch between c-MgZnO andYSZ, the c-MgZnO grown on YSZ shows the MgZnO(111) plane parallel to the(100) plane of YSZ. The off-plane XRD phi-scan of c-MgZnO shows 12-foldazimuth symmetry, implying multiple in-plane orientation of c-MgZnOlattice cells with respect to that of YSZ. Thus c-MgZnO can not beepitaxially grown on YSZ/Si(100). This challenge is overcome byinserting a thin layer of BTO (<100 Å) as the second buffer betweenc-MgZnO and YSZ. The high affinity of c-axis growth of BTO is helpfulfor c-axis growth of c-MgZnO. After the adoption of BTO, c-axisoriginated c-MgZnO films with 4-fold azimuth symmetry was realized onSi(100).

Other buffer layers such as CeO₂ and B₁₄T₁₃O₁₂ both have less than 1% oflattice mismatch to a Si(100) substrate. CeO₂ is also known forepitaxial growth on Si (111). However, when grown on Si(100), neitherone gives high quality epilayers, perhaps due to the oxidation of the Sisurfaces.

(b) Deviation of Zn and Mg composition in the c-MgZnO films from that inthe target—The energy bandgap of c-MgZnO is usually much larger thanthat of ZnO. To grow the c-MgZnO films, one may have surmised that a Mgrich PLD target is necessary. However, we have found that thecomposition of Mg in the films deviate significantly from the molepercentage of Mg in the target. Even a PLD target with a low Mgcomposition of 25% can be used for growing c-MgZnO films.

The unexpected higher Mg concentration in the films implies a net Znloss. Several potential causes include: 1) The MgZnO target was sinteredat 1300° C. and the Zn and Mg may evaporate, especially from the surfaceregion, in the sintering process with different ratios that cause a netZn loss; 2) during the PLD process, the Zn and Mg elements may havedifferent photo absorption cross-sections and thus different ablationrates; 3) when reached at the substrate, Zn species may also have asmaller sticking coefficient than that of Mg, thereby resulting in aMg-rich film; 4) the different vapor pressures of Zn and Mg also can notbe neglected; 5) in addition, Zn and Mg species may have differentoptimal windows of oxygen partial pressure-at high temperature, theactual partial oxygen pressure inside the PLD chamber may differ fromthe vacuum gauge reading due to the degassing of the heater, which mayalso cause different Zn and Mg deposition rates.

To unveil the cause of Zn lose, we first measured the Zn and Mgcomposition in the target using energy-dispersive-x-ray spectrometer. Nosignificant Zn loss was found from the target surface. Furthermeasurement of Zn and Mg depth profiles revealed that Zn and Mg haveuniform distributions in the target. Thus high temperature sintering wasruled out as the cause of Zn lose.

More clues come from the plot of Mg composition in the c-MgZnO films asa function of deposition temperature as shown in FIG. 7. Clearly thefilm deposited close to room temperature has almost the same Mg and Zncomposition as in the target, implying that the ablation rates of Mg andZn species is close if not equal. Thus, ablation rate difference shouldbe excluded as the cause of Zn loss in the films. This result alsoimplies that the sticking coefficient difference between Mg and Znspecies is minimal, at least for a substrate at room temperature. Wealso measured the Mg and Zn composition at higher growth temperaturesunder identical growth conditions but on different substrates whichinclude Si(100), YSZ(100), Al₂O₃(0001), SrTiO₃(100) and find similar Mgand Zn compositions. Consequently the suspicion of sticking coefficientcaused Zn loss should be excluded.

The influence of oxygen partial pressure on the changes of Mg to Znratio in the c-MgZnO films was also excluded, since the measuring of Mgto Zn ratio from the films deposited at the same temperature butdifferent oxygen pressure (10⁻¹ to 10⁻⁷ torr) have similar results.

After excluding other causes, the high vapor pressure of Zn spices atelevated temperature is most likely responsible for the Zn loss in thec-MgZnO films. At 400° C., the vapor pressure of Zn and Mg are 1×10⁻¹torr and 2×10⁻³ torr, respectively. These two values go up to 2×10² torrand 3×10¹ torr, respectively at 800° C. Qualitatively, such a high vaporpressure is enough to affect the Zn composition. When the vapor pressureof Zn and Mg is high enough, the Zn and Mg composition in the film isdetermined by the vapor pressure rather than the mol percentage of Mgand Zn in the target. Quantitative analysis further confirms the result.Assuming that the composition of Mg and Zn in the cMgZnO films isinversely proportional to the Zn and Mg's vapor pressure, the Zn and Mgratio in the c-MgZnO film can be easily calculated as 11% and 89%,respectively, which is very close, the Mg 87%, Zn 13% shown in FIG. 7.Further evidence comes from measuring the Zn composition in the filmsusing Mg_(x)Zn_(1−x)O targets with different x value. If the vaporpressure of Zn and Mg dominates the Zn and Mg composition, cMgZnO filmsfrom different targets should have the similar Zn composition at hightemperature, when the vapor pressure rather than the composition oftarget dominates the composition of the film. We found that thedifference of Zn composition in the targets was only reflected in thelow temperature region, where the Zn and Mg vapor pressure is too low tohave any significant effect. At 800° C., the Zn and Mg compositions inall of the films are Mg 87%, Zn 13% respectively, even though thedifferent targets of Mg_(0.3)Zn_(0.7)O, Mg_(0.4)Zn_(0.6)O andMg_(0.6)Zn_(0.4)O were used. Thus the high vapor pressure of Zn and Mg,as well as the ratio of vapor pressures is most likely the cause of thedeviation of Zn loss in the c-MgZnO films.

It is worthwhile to note that even at a temperature as high as 1000° C.,the vapor pressure of ZnO and MgO is in the order of 10⁻⁵ torr and 10⁻⁶torr, respectively, which should not affect the composition of Zn and Mgin the films. This result implies that in the laser plume Zn and Mgintermediate species exist as atoms and ions rather than as oxidemolecules, although they exist as oxide in the target as well as in thefilms.

1. A UV detector consisting essentially of a single active thin film ofa photoconductor material consisting of Mg_(x)Zn_(1−x)O, wherein x has avalue such that said photoconductor material will generate a detectableelectric current when exposed to UV light in the wavelength range offrom about 150 nm to about 400 nm, and first and second spaced apartelectrodes in contact with a top face of said active thin filmphotoconductor material.
 2. A visible-blind UV detector according toclaim
 1. 3. A solar-blind UV detector according to claim
 1. 4. A UVdetector according to claim 1, wherein x has a value between about 0and
 1. 5. A UV detector according to claim 1, wherein said materialcomprises Mg_(0.34)Zn_(0.66)O.
 6. A UV detector according to claim 1,including a substrate supporting a bottom face of said thin filmphotoconductor material.
 7. A UV detector according to claim 6, whereinsaid thin film is epitaxially grown on said substrate.
 8. A UV detectoraccording to claim 7, wherein said thin film is epitaxially grown byPLAD.
 9. A UV detector according to claim 6, including a buffer layerbetween said thin film and said substrate, said buffer layer comprisinga material that accommodates any lattice and/or thermal mismatch betweensaid thin film and said substrate.
 10. A UV detector according to claim9, wherein said buffer material comprises SrTiO₃/SrO, SrTiO₃/TiN orBiTiO₁₂/Y-stabilized ZrO₂.
 11. A UV detector according to claim 1,further comprising: a voltage source connected across said first andsecond electrodes, said voltage source creating an electric field withinsaid thin film, wherein when the surface of the thin film upon which theelectrodes are deposited is subjected to a photon illumination, electronhole pairs are created and flow within said thin film.
 12. A method offabricating a UV detector according to claim 11, comprising fabricatingsaid thin film, depositing a first electrode on a surface of said thinfilm, depositing a second electrode on the surface of the thin film,said second electrode being spaced from said first electrode, andconecting a voltage means and a current detector to said first andsecond electrodes.
 13. A UV detector according to claim 11, including acurrent detector connected to said first and second electrodes tomeasure the current from said thin film when exposed to UV radiation.14. A UV detector according to claim 11, wherein said voltage source isa battery.